High thermal stability and low Gilbert damping constant of CoFeB/MgO bilayer withperpendicular magnetic anisotropy by Al capping and rapid thermal annealingDing-Shuo Wang, Shu-Yu Lai, Tzu-Ying Lin, Cheng-Wei Chien, David Ellsworth, Liang-Wei Wang, Jung-Wei

Liao, Lei Lu, Yung-Hung Wang, Mingzhong Wu, and Chih-Huang Lai

Citation: Applied Physics Letters 104, 142402 (2014); doi: 10.1063/1.4870770 View online: http://dx.doi.org/10.1063/1.4870770 View Table of Contents: http://scitation.aip.org/content/aip/journal/apl/104/14?ver=pdfcov Published by the AIP Publishing

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High thermal stability and low Gilbert damping constant of CoFeB/MgObilayer with perpendicular magnetic anisotropy by Al capping and rapidthermal annealing

Ding-Shuo Wang,1 Shu-Yu Lai,1 Tzu-Ying Lin,1 Cheng-Wei Chien,2 David Ellsworth,3

Liang-Wei Wang,1 Jung-Wei Liao,1 Lei Lu,3 Yung-Hung Wang,2 Mingzhong Wu,3

and Chih-Huang Lai1,a)

1Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu, Taiwan2Electronics and Optoelectronics Research Laboratories, Industrial Technology Research Institute,Chutung, Taiwan3Department of Physics, Colorado State University, Fort Collins, Colorado 80523, USA

(Received 30 December 2013; accepted 27 March 2014; published online 7 April 2014)

We demonstrate that the magnetic anisotropy of the CoFeB/MgO bilayer can be manipulated by

adding an aluminum capping layer. After rapid thermal annealing, we can achieve large

perpendicular magnetic anisotropy of CoFeB with a high thermal stability factor (D¼ 72) while the

Gilbert damping constant can be reduced down to only 0.011 simultaneously. The boron and residual

oxygen in the bulk CoFeB layer are properly absorbed by the Al capping layer during annealing,

leading to the enhanced exchange stiffness and reduced damping. The interfacial Fe-O bonding can

be optimized by tuning annealing temperature and thickness of Al, resulting in enhanced

perpendicular anisotropy. VC 2014 AIP Publishing LLC. [http://dx.doi.org/10.1063/1.4870770]

Single CoFeB ultrathin film in CoFeB/MgO bilayer

structure was recently found to possess perpendicular mag-

netic anisotropy (PMA)1 and was recognized to be the most

promising candidate for high-performance spintronic devices

due to its large tunneling magnetoresistance (TMR) ratio2

among all PMA competitors. Many prototype devices have

been demonstrated by using this bilayer structure, such as

magnetic tunnel junctions (MTJs)1–8 and domain-wall-

motion based magnetic nanowires.9,10 However, relatively

low thermal stability11 and high Gilbert damping constant1

among those devices constrained this bilayer structure for

applications of high-performance spintronic devices.

Thermal stability factor D, defined as the ratio of static mag-

netic energy, KeffVact (Keff: effective anisotropy energy;

Vact: activation volume) to thermal energy, kBT (kB:

Boltzmann constant), is one important indicator for informa-

tion retention against thermal fluctuation. However, the

reported D of a single p-CoFeB layer is still insufficient for

stable information storage. The PMA in CoFeB/MgO bilayer

is believed to originate from the interfacial anisotropy at

interface12–14 and can be manipulated by interface engineer-

ing techniques.14,15 To establish high D in p-CoFeB, double

CoFeB layers with a Ta insertion layer were reported.16 The

increased magnetic volume has indeed enhanced D to �100,

but the effects of Ta insertion on the damping constant and

magnetization reversal are unclear.

Furthermore, the critical switching current density Jc0

for spin-transfer torque (STT) devices is proportional to

Gilbert damping constant a of the free layer.17 It was

reported that a is up to 0.027 for 1.3 nm CoFeB in p-MTJ1

and would be increased with further decreased thickness.

Other works focusing on lowering the Gilbert damping

constant of p-CoFeB reported that a could be tuned by opti-

mizing composition and crystallinity,18 or by inserting a

metal Mg layer at CoFeB/MgO interface.15 In this work, we

report an alternative way to effectively modify the magnetic

properties of CoFeB/MgO by introducing an Al capping

layer combining with post rapid thermal annealing (RTA).

We achieve both high thermal stability and low Gilbert

damping constant simultaneously, which makes the bilayer

structure more suitable for device applications.

Film stacks of substrate // Pd (3)| Co20Fe60B20 (target

composition) (1)| MgO (1.6)| Al (x) (in nanometers), where

x was 0 to 5 nm, were deposited on thermally oxidized silicon

substrates by sputtering under a base pressure <5� 10�8 Torr.

The MgO layer was deposited by RF sputtering from an MgO

target while other layers were deposited by DC sputtering. We

deposited the MgO layer by two separated steps. After the dep-

osition of first MgO layer (0.4 nm), an oxygen plasma treat-

ment was applied to the first MgO layer to adjust the amount

of excess oxygen at CoFeB/MgO interface. The second MgO

layer (1.2 nm) was deposited subsequently onto first MgO

layer. The film stack was annealed in a RTA furnace in vac-

uum at elevated temperatures for 30 s without external mag-

netic field. The film thickness was calibrated by atomic force

microscopy (AFM) and cross-sectional transmission electron

microscopy (XTEM) images. The static magnetic properties,

saturation magnetization Ms and anisotropy field Hk, were

measured by vibrating sample magnetometer (VSM), where

Hk was obtained from the saturation field along the hard axis.

The thermal stability factor D was deduced by using waiting

time method,19 a commonly used technique in magnetic re-

cording media. After positive saturation, we applied a negative

field H (<Hc) and the characteristic relaxation time t50 of mag-

netic moments can be expressed by

t50 ¼ eaðH��HÞ; (1)

a)Author to whom correspondence should be addressed. Electronic mail:

chlai@mx.nthu.edu.tw

0003-6951/2014/104(14)/142402/5/$30.00 VC 2014 AIP Publishing LLC104, 142402-1

APPLIED PHYSICS LETTERS 104, 142402 (2014)

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where a¼MsVact/kBT and H* are fitting parameters. Vact

represents the activation volume of the reversed domain, and

D can be further calculated by

D ¼ E

kBT¼ Kef f Vact

kBT; (2)

where E is the energy barrier height for magnetization rever-

sal, and Keff is effective anisotropy energy equal to MsHk/2.

The Gilbert damping constant a was resolved from fer-

romagnetic resonance (FMR) spectra acquired by placing the

film stack with 1 cm2 on a 50 X co-planar waveguide (CPW)

structure that consisted of a 50-lm-wide signal line and two

grounds 25 lm away from signal line with the substrate side

facing up, and measuring the transmission coefficients of

the film/CPW structure with a vector network analyzer

(VNA).20,21 The sample was coated with 1 lm photoresist to

prevent shorting of the ground and signal lines. In our exper-

imental setup, we first fixed the input microwave frequency

and swept the magnetic field to measure one FMR spectrum

followed by repeating the field sweep with various micro-

wave frequencies over a range of 5–33 GHz. The Kittel for-

mula is used to express the FMR frequency

fFMR ¼c

2p

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiHFMRðHFMR þ 4pMef f Þ

q; (3)

where HFMR is the applied field at resonance, 4pMeff repre-

sents the effective magnetic anisotropy field, and c¼ glB/�is the absolute gyromagnetic ratio (g is the Land�e g-factor).

Furthermore, a and the inhomogeneous linewidth broadening

DHILB can be extracted from the FMR peak-to-peak line-

width DH by20,21

DH ¼ 2ac

2pfFMR þ DHILB: (4)

The chemical states and element distribution in the mul-

tilayer structure were probed by X-ray photoemission spec-

troscopy (XPS) equipped with an Ar ion-milling gun for

resolving depth profile, performed at BL24A of NSRRC,

Taiwan. The best energy resolution of the energy analyzer

was 0.02 eV. The energy of incident photons was calibrated

by metallic Au 4f7/2 peak at 84 eV from a standard clean

gold chip.

Fig. 1(a) shows the in-plane (IP) and out-of-plane (OOP)

hysteresis loops of the as-deposited CoFeB/MgO sample. It

exhibits in-plane magnetic anisotropy (IMA) with Hk (OOP

saturation field Hs,?) of �3000 Oe. After 360 �C annealing,

the non-capped sample presents low PMA with small OOP re-

manent magnetization and perpendicular coercivity Hc,?(shown in Fig. 1(b)), whereas its Hk is �1200 Oe. Ms is

slightly increased from 200 emu/cc to �250 emu/cc after

RTA. Interestingly, as shown in Fig. 1(c), the 5 nm-Al-capped

sample exhibits much stronger PMA after same RTA process;

its Hc,? is �500 Oe and Hk> 4000 Oe. In addition, Ms signifi-

cantly increased to �1500 emu/cc after RTA. Figs. 1(d) and

1(e) summarize the RTA temperature dependence of Ms and

Hk for both 5 nm-Al-capped and non-capped samples, respec-

tively. For Al-capped samples, Ms grows drastically with

increasing annealing temperature (Tann) and saturates at

360 �C. However, for non-capped samples, the Ms increment

is relatively insignificant with a maximum only �300 emu/cc

after 450 �C RTA. Hk is substantially increased with Tann and

reaches a maximum �4200 Oe in Al-capped samples.

Moreover, the direction of magnetic easy axis dramatically

changes from OOP to IP when Tann exceeds 400 �C. For

non-capped samples, however, the Hk is small and insensitive

to temperature when Tann is above 270 �C. We further varied

the Al thickness and observed that the PMA-IMA transition

temperature (Ttrans) was increased linearly by reducing Al

thickness down to 2 nm, as shown in Fig. 1(f). We also adjust

the oxygen plasma treatment time tox and found out that Ttrans

is strongly correlated and increased with tox (data not shown).

For spintronic devices, thermal stability factor D and

Gilbert damping constant a are two critical properties. We

selected 5 nm-Al-capped samples with RTA temperature

ranging from 330 �C to 390 �C, which revealed good PMA

with large remanent magnetization (Mr/Ms> 0.9), to analyze

D. The waiting time dependences on the applied field are

summarized in Fig. 2(a). The nucleation diameter Unucl and

D can be calculated from the fitted Vact and Eq. (2), respec-

tively. Here, we assumed cylindrical shaped nucleation

so that Unucl is equal to 2(Vact/pt)1/2, where t is the CoFeB

thickness. Note that we analyze the unpatterned ultrathin

film whose thickness (1 nm) does not allow thermal magnons

to transport along OOP direction, and the lateral

dimension of the film is much larger than the exchange

length kex� (Aex/Keff)1/2¼ (2Aex/MsHk)1/2, where Aex is the

exchange stiffness constant; therefore, the magnetization re-

versal should proceed through sub-volume excitation rather

than macrospin-like.22 The deduced Unucl from waiting time

measurement is about 30–50 nm (Eq. (1)), similar to the

order of exchange length kex in CoFeB.22 As shown in Fig.

2(b), Keff increases monotonically whereas Unucl behaves

oppositely with Tann. However, the product of Keff and Vact,

representing energy barrier height E for reversal, is nearly

unchanged with Tann within our experiment regime, as

FIG. 1. IP and OOP hysteresis loops of CoFeB/MgO bilayer with different

post-treatments: (a) as-deposited, (b) 360 �C RTA for 30 s without Al cap-

ping, and (c) 360 �C RTA for 30 s with 5 nm Al capping. Insets of (a) and

(b) show the enlarged loops near zero magnetic field. Dependences of (d) Ms

and (e) Hk on RTA temperatures for non-capped and Al-capped samples

(�Hk: IMA). (f) Variations of PMA-IMA transition temperature with Al

capping thickness.

142402-2 Wang et al. Appl. Phys. Lett. 104, 142402 (2014)

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shown in Fig. 2(c). In the scenario of sub-volume excita-

tion,22 E is of the order �4pAext, only controlled by Aex and

t; but in our case, E is proportional to Aex only since t is

unchanged. Our result shows a high D (�72), greater than

most of previously reported results for a single p-CoFeB

layer,11 implying strong exchange stiffness in our samples.

The calculated D is almost unchanged since Aex may not

vary significantly with temperature in our experimental re-

gime. Based on the sub-volume excitation mode,22 Vact

would be inversely proportional to Keff, which is consistent

with our observation, as shown in Fig. 2(b). It seems that we

did not really gain advantages for promoting D by increasing

Keff; however, for high density devices, the films need to be

patterned into a small size with a lateral dimension less than

kex; in this case, the volume associated with D would be the

physical size of patterns; therefore, D can be enhanced by

increased Keff.

Fig. 3(a) shows the dependence of FMR frequency on

magnetic fields for 360 �C RTA (with PMA) and 450 �CRTA (with IMA) samples, respectively. We calculate the

effective magnetic anisotropy field 4pMeff by using Eq. (3)

and Fig. 3(a). The 4pMeff obtained by using VNA-FMR

coincided with Hk measured by using VSM quite well in our

experiments. Fig. 3(b) shows the dependence of FMR peak-

to-peak linewidth DH on microwave frequency; a was

deduced by using Eq. (4). The fitted a of the 360 �C RTA

sample (PMA) is 0.011 6 0.002, while it is 0.016 6 0.001 for

the 450 �C RTA sample (IMA). The a of the PMA sample is

quite low comparing with previously reported values of the

single CoFeB layer with PMA.1,15,18

To understand the role of Al capping layer, XPS experi-

ments were carried out to investigate the chemical states and

element distributions in the multilayer structure of samples

with different post-treatments. Fig. 4(a) shows the Fe2pspectra of 5 nm-Al-capped samples after different post-

treatments. The satellite peaks in Fe spectrum of the as-

deposited sample indicated that Fe was oxidized, which

could originate from RF-sputtering deposition of adjacent

MgO layer and oxygen plasma treatment. After 360 �C RTA,

the satellite peaks disappeared, and only a small shoulder

left near Fe 2p3/2 in spectrum (indicated by an arrow), indi-

cating that the amount of Fe oxides was significantly dimin-

ished with limited Fe-O bonding left. Majority of Fe was

reduced to be metallic after 450 �C RTA. The spectra of Co

revealed that Co was oxidized in the as-deposited sample but

were reduced to be metallic after 360 �C RTA (data not

shown). Since the PMA in CoFeB/MgO system resulted

from the interfacial Fe-O orbital hybridization, the signifi-

cant reduction of Fe oxides may substantially lower the

PMA, leading to the PMA-IMA transition at 400 �C as

shown in Fig. 1(e). In the non-capped sample, however, both

Fe and Co mostly remained oxidized even after 450 �C RTA

(data not shown).

In addition, the Al2p and B1s X-ray photoemission

spectra taken at different ion-milling stages are shown in

Figs. 4(c) and 4(d) of as-deposited and 360 �C RTA 5 nm-

Al-capped samples, respectively. The series of Al2p spectra

of the as-deposited sample revealed that a small amount of

AlOx were observed at MgO/Al interface, which may result

from chemisorbed oxygen species or the partial intermixing

during deposition, while the rest regions of Al were still me-

tallic. However, after 360 �C RTA, Al has mostly been oxi-

dized, and the signal of AlOx was even increased for the

regions away from MgO/Al interface, indicating that O

atoms have diffused out from the bottom part of the sample

FIG. 2. (a) Waiting time dependence on reversed external magnetic field of

Al-capped samples with various RTA temperatures; the solid lines represent

the fitting results of Eq. (1). Variations of (b) nucleation diameter Unucl and

effective anisotropy energy Keff and (c) thermal stability factor D with

annealing temperature Tann.

FIG. 3. (a) FMR frequency dependence on external field and (b) FMR peak-

to-peak linewidth dependence on microwave frequency of 360 �C RTA and

450 �C RTA Al-capped samples.

FIG. 4. (a) Fe2p X-ray photoemission spectra of Al-capped samples with

different post treatments: as-deposited, 360 �C RTA and 450 �C RTA,

respectively. “Sat.” represents satellite peaks of Fe oxide in the as-deposited

sample. (b) Schematic illustration of X-ray probing regions at various sput-

tering stages. B1s and Al2p X-ray photoemission spectra taken at different

sputtering stages of (c) as-deposited and (d) 360 �C RTA Al-capped samples,

respectively. The probing regions in each curve are shown in (b) by different

symbols.

142402-3 Wang et al. Appl. Phys. Lett. 104, 142402 (2014)

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to the whole Al layer. The base pressure of our RTA system

is <1� 10�5 Torr, and the amount of residual O in chamber

would not be sufficient to cause the observed changes (Fig.

4(d)) during RTA. To exclude the oxygen effects during

RTA, another 5 nm Al film with a Pd buffer layer deposited

on the same substrate and RTA at 360 �C is used as a refer-

ence sample. We do not observe distinct difference in XPS

spectra of Al2p for the as-deposited and 360 �C RTA sam-

ples, which reconfirms the observed changes of O distribu-

tions are caused by the O diffusion from the MgO/Al

interface. The series of B1s spectra, shown in Fig. 4(c), show

that the BOx signal can only be observed near MgO/Al inter-

face in the as-deposited sample, which could be due to finite

XPS probing depth, that is, the majority of B atoms locate in

the CoFeB layer in the as-deposited stage. Oxidation of B in

CoFeB may be related to our oxygen plasma treatment. After

360 �C RTA, a clear BOx signal appears on the top surface

of the Al layer, as shown in Fig. 4(d), which implies that lots

of B atoms have diffused out from CoFeB to Al layer. The

XPS result revealed that B and O atoms diffused from the

bottom CoFeB to top Al capping layer and reduction of Co

and Fe occurred accordingly during RTA. The evolution of

magnetic properties with Tann can be understood as a result

of B and O diffusion as well as the reduction of magnetic

metals. Al2O3 is a more stable oxide than Fe and Co oxides;

furthermore, B and Al also have strong affinity. Therefore,

the Al capping layer may behave as the O and B sink to drive

both species to diffuse out of CoFeB layer. Thicker Al layer

might provide higher driving force for diffusion, therefore,

the onset temperature for significant reduction of magnetic

metals, which results in the PMA-IMA transition, is lower,

as we observed in Fig. 1(f). The elimination of B and O in

the bulk CoFeB layer for the sample with Al capping after

RTA leads to strong Aex and hence high D. Due to the short

period of RTA process, we may maintain the required Fe-O

bonding at the interface, which give rise to high perpendicu-

lar anisotropy. On the contrary, the non-capped sample does

not have sufficient driving force to re-distribute B and O;

therefore, the change of magnetic properties through RTA is

limited.

In addition to the compositional changes, the structures

with different RTA temperatures are also varied. Fig. 5 shows

the XTEM images of Al-capped sample with different post

treatments: (a) 360 �C RTA and (b) 450 �C RTA, respec-

tively. The main structural difference between two samples is

the crystallinity of CoFeB. The 450 �C RTA sample reveals a

well-crystallized structure but the 360 �C RTA sample

remains mostly amorphous in the bulk region away from

CoFeB/MgO interface. Since intrinsic a comes from spin-

lattice scattering, the improved crystallinity for high-

temperature annealing samples leads to a higher a. Our

results suggest that the reduced a by combining Al capping

layer with RTA process can be attributed to the eliminated B

content and suppressed re-crystallization of bulk CoFeB. We

speculate that the low temperature RTA process can only

trigger the diffusion of light atoms (B, O) to be absorbed by

Al capping layer and reduces the oxides simultaneously, but

does not provide sufficient energy for crystallization of bulk

CoFeB. That is, after RTA, we can eliminate B and excess O

atoms to achieve high Aex and PMA, which result in

enhanced D; meanwhile, the bulk crystallinity is remained

relatively low so a stays low. For higher temperature RTA,

however, the thermal energy is high enough for significant

reduction of interfacial oxides and to promote crystallinity of

CoFeB, yielding to observed PMA-IMA transition and

increased a. As a comparison, long annealing time in conven-

tional annealing system can also enhance the crystallinity of

300 �C-annealed CoFeB, leading to a larger a.18 Furthermore,

we found out that the CoFeB layer started to crystallize from

the CoFeB/MgO interface since the Al capping may proc-

esses stronger affinity for B than Pd.23 The orientation of

as-crystallized CoFeB followed the MgO (002) orientation,

however, at higher annealing temperatures, part of CoFeB

layer also crystallized from the bottom Pd/CoFeB interface

and followed the Pd (111) orientation.24 The influence of

crystal orientation on PMA in p-CoFeB is unclear. On the

other hand, when we change the initial oxidation period, we

do observe a strong correlation between initial amounts of ox-

ygen on PMA-IMA transition but we do not observe the

structural differences for the samples with various initial oxi-

dation period; therefore, we believe that the reduction of

Fe-O bonding is the dominant factor to lower the interfacial

anisotropy, resulting in the observed PMA-IMA transition.

Because all excess O and B atoms were removed out of

CoFeB bulk, our approach has also raised the Ms of CoFeB,

which may increase Jc0 in STT devices; however, low awould partially compensate this unfavorable factor with

gained high D. Furthermore, in recent spintronic device

designs using spin Hall effect or inversed spin Hall effect

through spin-pumping, Ms is proportional to the spin-mixing

conductance (g"#) and the output electromotive forces;25

therefore, high Ms may result in large output signals for these

devices. The modification and optimization of interfacial and

bulk composition as well as structure in CoFeB/MgO struc-

ture are quite challenging by simply tuning deposition condi-

tions. Our method, in fact, provides an alternative solution to

enhance the required magnetic properties of this bilayer

structure.

In summary, we manipulate the magnetic properties of

p-CoFeB by introducing an additional Al capping layer with

a RTA process, and D can be enhanced up to �72 while acan be reduced down to �0.011 simultaneously.

This work was financially sponsored by National

Science Council of Taiwan, Republic of China under Grant

No NSC-100-2221-E-007-064-MY2. The authors would like

FIG. 5. TEM cross-sectional images of 5 nm-Al-capped sample after differ-

ent post treatments: (a) 360 �C RTA and (b) 450 �C RTA.

142402-4 Wang et al. Appl. Phys. Lett. 104, 142402 (2014)

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123.195.21.113 On: Mon, 07 Apr 2014 13:55:18

to thank Professor Yaw-Wen Yang, Dr. Liang-Jen Fan, and

Dr. Chia-Hsin Wang at National Synchrotron Radiation

Research Center (NSRRC) for their support on XPS experi-

ments. The HRTEM analysis is partly supported by Mr.

Hsueh-Ren Chen at Instrumentation Center of National

Taiwan University.

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